Carbide grain growth in cemented carbides sintered with alternative binders

Carbide grain growth in cemented carbides sintered with alternative binders

Journal Pre-proof Carbide grain growth in cemented carbides sintered with alternative binders Z. Roulon, J.M. Missiaen, S. Lay PII: S0263-4368(19)30...

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Journal Pre-proof Carbide grain growth in cemented carbides sintered with alternative binders

Z. Roulon, J.M. Missiaen, S. Lay PII:

S0263-4368(19)30466-4

DOI:

https://doi.org/10.1016/j.ijrmhm.2019.105088

Reference:

RMHM 105088

To appear in:

International Journal of Refractory Metals and Hard Materials

Received date:

14 June 2019

Revised date:

3 September 2019

Accepted date:

10 September 2019

Please cite this article as: Z. Roulon, J.M. Missiaen and S. Lay, Carbide grain growth in cemented carbides sintered with alternative binders, International Journal of Refractory Metals and Hard Materials(2018), https://doi.org/10.1016/j.ijrmhm.2019.105088

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© 2018 Published by Elsevier.

Journal Pre-proof

Carbide grain growth in cemented carbides sintered with alternative binders Z.Roulon*, J.M Missiaen, S.Lay Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMAP, F-38000 Grenoble, France * [email protected]

Abstract

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Cemented carbides are composites made of a hard refractory ceramic phase and a ductile binder, most commonly WC and Co, respectively. Since the use of cobalt in the hard metal industry is questioned by the new European regulation on chemicals, extensive research has been done to develop new grades based on a Co-Ni-Fe binder. With similar mechanical, physical properties and affinity to C and W, nickel and iron are the best candidates for an efficient binder in cemented carbides. As mechanical properties are strongly dependent on the materials microstructure, and especially on the WC grain size, understanding the effect of the binder on the final microstructure is crucial. In this work, the carbide grain growth behaviour of WC-M alloys (M=Co, Ni, Fe) with different carbon contents is discussed from qualitative and quantitative microstructural analyses. Whereas grain growth is more or less inhibited in WC-Fe alloys, increasing carbon content promotes grain growth in WC-Co and WC-Ni alloys, with a slight abnormal grain growth in case of Ni binder. Different mechanisms for grain growth are discussed, in relation with the observed morphology of WC grains after sintering.

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Keywords: cemented carbides, sintering, alternative binder, grain growth, microstructure

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Journal Pre-proof 1. Introduction

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Cemented carbides have been used in the cutting tools industry since the 1920’s as substitute to high costly diamond tools [1].These materials are usually processed by the conventional powder metallurgy route [2]. Known as a very hard ceramic material, tungsten carbide is used as the major component of cemented carbides. A ductile metal binder is added as the matrix, which in most cases is cobalt [3]. Co is the most versatile binder because it provides excellent wetting on the tungsten carbide grains, and has superior mechanical properties than other metallic alloys [4]. Moreover, the combination of hardness and ductility provides exceptional properties to the cemented carbide alloys. Nevertheless, the use of cobalt as a binder is questioned by the new European regulation on chemicals [5]. Therefore, new alternative binders are considered, especially Fe and Ni-alloys [6][7]. Both iron and nickel metals have been investigated as candidates, since they are the closest transition metals from cobalt in the periodic table and by consequence have a similar affinity with carbon and tungsten [8], in addition to similar mechanical properties.

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Control of the final microstructure is critical to reach high final mechanical properties. Since the beginning of the cemented carbides industry, researches have been conducted to understand grain growth mechanisms involved in cemented carbides during sintering stage.

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Grain growth of solid particles in a liquid is usually described by Ostwald ripening i.e the dissolution of the smallest grains, followed by diffusion of solute species in the liquid and uniform precipitation on bigger grains. LSW theory (Lifshitz and Slyozov [9] , and Wagner [10]) modelled grain growth kinetics in stationary conditions, assuming a population of spherical particles dispersed in a liquid phase. The limiting process can be either solute diffusion or dissolution/precipitation reaction. The particle size distribution remains selfsimilar with time according to the theory. In the case of WC-Co materials, evolution of the grain size calculated for a diffusion-controlled grain growth is significantly faster than observed experimentally, which suggests that diffusion in the liquid is not the limiting step. In addition, the particle size distribution is significantly broader than the theoretical LSWpredicted one [11] [12] and evolution with time is not self-similar [13], which suggests that precipitation at the particle surface of the grains is not uniform but occurs on preferential sites. For faceted grains, as WC in WC-Co alloys, an energy barrier is needed for the nucleation of the species at the surface or on existing defects of the grains [11]. Therefore, grain growth may be limited in this case by 2D-nucleation or nucleation on defects, or by the lateral growth of nuclei at the interface. The driving force for grain growth depends on interface energy and effective radius of the grains. Following Kang [14], abnormal grain growth occurs as the driving force of the bigger grain is superior to a critical driving force. At last, WC/WC grain boundaries have to be taken into account, since WC particles are not surrounded by the liquid. Grain growth needs a cooperative migration of phase boundaries and grain boundaries [15][16]. Grain boundaries and phase boundaries can either impede or enhance grain growth [17], depending on their relative mobility and on the fractional area of grain boundaries on the particle surface, i.e the contiguity [16] [18]. In addition, grain growth may be facilitated for specific local configurations with low energy or high mobility grain boundaries and this may be an additional reason for the broadening of the particle size distributions, as compared to predictions of the LSW theory [15]. A lot of work has been performed on the effect of grain size distribution [11] [19][20], carbon balance [21][22][23] and grain growth inhibitors [24] on WC grain coarsening in WC-Co cemented carbides but very few with alternative Fe or Ni binder. Considering the effect of the binder nature, Wittmann [25] observed a more important grain growth with a Ni binder than 2

Journal Pre-proof with a Co binder while grain growth was strongly restricted with a Fe binder, but no quantitative analysis has been done. Moreover, more detailed observations of the WC/binder interfaces are required to specify grain growth mechanisms in relation with the different binders. In this work, the WC grain growth behavior of WC-M alloys (M=Co, Ni, Fe) is studied from quantitative microstructural analyses as a function of sintering time at 1450°C. Quantitative analysis is based on intercept length distributions obtained by image analysis on the WC and binder phase. Finally, WC grain growth mechanisms are discussed from the experimental results and the literature.

2. Experimental conditions

WC-Ni C-rich 5.49

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WC-Co W-rich 5.22

WC-Ni W-rich 4.98

WC-Fe C-rich 5.93

WC-Fe W-rich 5.30

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C wt [%]

WC-Co C-rich 5.60

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Tab.1: Carbon content in wt [%] of the WC-M (M=Co,Ni,Fe) alloys

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Six alloys with approximately 20 vol% binder and different carbon contents were investigated (Tab.1). The C-rich alloys are in the 3-phase (WC-M-C) domain and the W-rich ones in the (WC-M-M6C) domain as indicated by dashed lines on Fig. 1, Fig. 2 and Fig. 3.

Fig. 1: Equilibrium phase diagram of WC-20vol%Co (TCE_09 database). Dashed red lines indicate the carbon composition of the alloys (Tab.1).

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Fig. 2: Equilibrium phase diagram of WC-20vol%Ni (TCE_09 database). Dashed red lines indicate the carbon composition of the alloys (Tab.1).

Fig. 3: Equilibrium phase diagram of WC-20vol%Fe (TCE_09 database). Dashed red lines indicate the carbon composition of the alloys (Tab.1).

Because of the narrowness of the 2-phase region of WC-Fe and the tendency of the alloys to loose carbon during the sintering process [26], choice has been made to investigate 3-phase alloys on both sides of the carbon window for the 3 binders.

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Alloys are made of pre-alloyed WC powder and raw metallic powder (Fe, Ni, Co), with a mean size from 2 to 8 µm approximately. Milling step was proceeded by ball milling with cylindrical milling bodies, for approximately 43 h in ethanol. Then the mills were spray-dried to remove the alcohol and granulated in final particles with approximately 5 to 30 μm diameter size. Samples were compacted in a 8 mm diameter cylindrical die to reach 55% green density. After debinding in He-H2 atmosphere, sintering was performed in a SETARAM TMA92 dilatometer in Ar atmosphere, with a heating rate of 3°C/min, a sintering plateau at 1450°C for 0.5h, 4h or 8h and a cooling rate of 30°C/min. Scanning Electron Microscope (SEM) experiments were conducted using a SEM-FEG Zeiss Ultra 55. The sintered samples were cut along a vertical section by a diamond wire and polished with grinding steps from 40 µm to 1 µm. SEM images were acquired in the backscattered mode. The accelerating voltage was fixed at 10 kV to limit the penetration depth of the electrons. Ten 2048 x 1536 pixels pictures were taken at x3000 magnification by uniform sampling on the polished sections. The WC phase, binder phase and porosity, were segmented on SEM images and intercept length distributions were determined by image analysis techniques with Aphelion software ™. The mean intercept length was calculated together with the ratio between the intercept length in the 9th to the 5th decile of the distribution, l90/l50, to quantify a possible abnormal grain growth.

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Morphology of the WC grains was also observed by SEM, after removal of the binder from sintered material by heating for 24h in a mixture of hydrochloric acid and water [27]. In order to investigate the WC/binder interfaces and WC/WC grain boundaries, TEM experiments were performed on thinned specimens of material, using a JEOL 4000EX microscope.

3. Experimental results

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3.1. Microstructure evolution during isothermal sintering

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Microstructure evolution was characterized by SEM images ( Fig. 4 and Fig. 6) and image analysis (Fig. 5, Fig. 7 and Fig. 8) for WC-Fe, WC-Ni and WC-Co C-rich and W-rich sintered materials in order to study the influence of carbon content and sintering time on WC grain growth. In the case of W-rich alloys, grain growth is limited with the Co or Ni binder and practically inexistent with the Fe binder (Fig. 4, Fig. 7 and Fig. 8). Regarding the phase distribution (Fig. 4), η-phase is homogeneously distributed in WC-Fe Wrich material whereas it is heterogeneously distributed in WC-Co and WC-Ni W-rich alloys. η-phase is identified as light grey grains, with almost the same contrast than WC grains.

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Journal Pre-proof WC-Ni W-rich

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Fig. 4: Microstructures of WC-Ni W-rich, WC-Co W-rich and WC-Fe W-rich sintered at 1450°C for 0.5h, 4h and 8h

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Fig. 5: Intercept length distribution in the WC phase in the alloys a) WC-Ni W-rich, b) WC-Co W-rich, and c) WC-Fe W-rich at 0.5h, 4h and 8h of sintering at 1450°C.

Considering the C-rich alloys, influence of the binder nature on WC grain growth is different. Indeed, there is a significant grain growth for WC-Ni alloy (Fig. 6 and Fig. 7), with a major increase of mean intercept length and a significant widening of the distribution as sintering time is increased (Fig. 8). WC-Co presents a moderate grain growth from 0.5h to 4h sintering with little change from 4h to 8h (Fig. 6, Fig. 7 and Fig. 8). Unlike WC-Co and WC-Ni alloys, WC-Fe does not show any significant grain growth, with almost no change in the distribution of intercept length. After 8h sintering, WC-Ni C-rich has the coarser microstructure, followed by WC-Co C-rich then WC-Fe C-rich.

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Journal Pre-proof WC-Ni C-rich

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Fig. 6: Microstructures of WC-Ni C-rich, WC-Co C-rich and WC-Fe C-rich sintered at 1450°C for 0.5h, 4h and 8h

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Fig. 7: Intercept length distribution in the WC phase in the alloys a) WC-Ni C-rich, b) WC-Co C-rich, and c) WC-Fe C-rich at 0.5h, 4h and 8h of sintering at 1450°C.

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Journal Pre-proof WC-Ni 3 phase C-rich WC-Ni 3 phase W-rich 12

b) 3.2

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WC-Co 3 phase C-rich WC-Co 3 phase W-rich

l90/l50

Mean intercept length [µm]

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WC-Fe 3 phase C-rich WC-Fe 3 phase W-rich

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Fig. 8: a) Mean intercept length and b) l90/l50 as a function of sintering time

3.2. WC morphology and interfaces in WC-M alloys

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Whereas quantitative information about size distribution in the WC phase can be extracted from 2D image analysis, it is more difficult to evaluate the morphology of the grains with this technique. Observation of the microstructure after dissolution of the binder provides qualitative information about grain shapes and interface roughness. Steps at the interfaces between WC and binder grains are found in W-rich alloys, whatever the binder (Fig. 9). This is confirmed by TEM observations (Fig. 10), with a step size around 10 to 30 nm. They are observed on the basal planes of WC grains, particularly in the case of iron binder (Fig. 9.c, Fig. 10.c).

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Fig. 9: WC grains morphology from a) WC-Ni, b) WC-Co and c) WC-Fe W-rich alloys sintered at 1450°C for 30 min after chemical etching

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Fig. 10: Interfaces between WC and binder grains showing steps in a) WC-Ni, b) WC-Co, c) WC-Fe W-rich alloys

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In the case of C-rich alloys, WC grains have smooth facets without any steps and the global WC grain shape is a more or less a truncated triangular prism, with more rounded edges between basal and prismatic facets for the WC-Ni C-rich alloy (Fig. 11 & Fig. 12). Whatever the binder, no preferential growth direction is identified from these microscopic observations.

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Fig. 12: Interfaces between WC and binder grains in a) WC-Ni, b) WC-Co, c) WC-Fe C-rich alloys

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Fig. 11: WC grains morphology from a) WC-Ni, b) WC-Co and c) WC-Fe C-rich alloys sintered at 1450°C for 30 min after chemical etching

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Journal Pre-proof 4. Discussion Both 2D and 3D microstructure analyses provide information about WC grain growth kinetics and mechanisms as a function of binder nature and carbon content. Carbon content has a strong influence on the grain growth of cemented carbides, as already observed with WC-Co alloys [16] [22] [23]. A significant WC grain growth is observed for Crich alloys with Co and even more with Ni binder, whereas the microstructure evolution is much slower in the case of W-rich alloys. No significant grain growth is noticed for WC-Fe compositions, whatever the C content. These results confirm the qualitative observations of Wittmann [25].

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Regarding the WC morphology after sintering, the presence of steps exhibits the influence of carbon content on grain growth mechanisms occurring in the alloys. With carbon content below the stoichiometric WC, triangular terraces are observed on the basal planes, suggesting that 2D nucleation of new layers is much more rapid than the lateral growth of these layers. On the other hand, the absence of steps suggests than lateral growth does not limit the kinetics for C-rich alloys. In the case of WC-Ni C-rich alloys, the more rounded grain corners suggest than 2D nucleation and lateral growth do not limit grain growth kinetics anymore but that precipitation tends to become uniform at the particle interface. It can then be assumed that precipitation first occurs on rounded edge surface followed by atomic redistribution by fast surface diffusion on flat surface (Fig. 13).

Fig. 13: Schematic illustrating grain growth by 1 bulk .diffusion of species through liquid phase then 2.uniform precipitation on rounded WC edges, followed by 3.surface diffusion

The absence of energy barrier for the precipitation would also explain why the grain growth rate is larger for WC-Ni compared to WC-Co and WC-Fe C-rich alloys (Fig. 8). Indeed, from Fig. 8, grain growth rate can be estimated using intercept length measurement, and it is larger by roughly one order of magnitude for WC-Ni compared to WC-Co and WC-Fe C-rich alloys (approximately 1 µm/h, 0.1 µm/h and <0.1 µm/h respectively).

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Journal Pre-proof 5. Conclusion Grain growth in cemented carbides with a Co, Ni and Fe rich binder has been analyzed from a systematic study of intercept length distributions in the WC phase. Whereas grain growth is practically inhibited with Fe binder for all carbon contents, grain growth is rather slow for Ni and Co binder with low carbon content but it is enhanced and the abnormal character of grain growth increases when the carbon content increases, especially in the case of the Ni binder.

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WC/binder interface morphology gives information about grain growth mechanisms involved in the cemented carbides. Steps at the WC/binder interface suggest that lateral growth limits the kinetics in W-rich alloys whereas smooth interfaces indicate that this is no more the case in C-rich alloys. In addition, for WC-Ni C-rich alloys, apparition of more rounded grains suggest a more uniform precipitation of the solid on the rounded edges of WC grains.

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Acknowledgments The authors are grateful to Sandvik Coromant R&D, Stockholm, Sandvik Mining and Rock Technology, R&D Rock Tools, Stockholm and Seco Tools AB, Fagersta, Sweden for their financial and technical support, for provision of powder mixtures and for fruitful discussions about experimental results.

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Graphical abstract Highlights  Comparison of WC grain growth with Co, Ni and Fe binders in cemented carbides  Quantitative analysis from intercept length measurements  Discussion of mechanisms based on thorough microstructural observations  Precipitation facilitated on rounded grain edges for WC-Ni C-rich alloys

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